Assessment of Retained Austenite in Fine Grained Inductive Heat Treated Spring Steel

materials Article Assessment of Retained Austenite in Fine Grained Inductive Heat Treated Spring Steel Anna Olina 1,*, Miroslav Píška 1, Martin Petrenec 1, Charles Hervoches 2, Pˇremysl Beran 2,3, JiˇríPechoušek 4and Petr Král5 1Department of Manufacturing Technology, Faculty of Mechanical Engineering, Brno University of Technology, Technicka 2, 61669 Brno, Czech Republic; [email protected].cz (M.P.); [email protected] (M.P.) 2Nuclear Physics Institute of the CAS, ˇ Rež 130, 25068 ˇ Rež, Czech Republic; [email protected] (C.H.); [email protected] (P.B.) 3European Spallation Source ERIC, Box 176, 22100 Lund, Sweden 4Department of Experimental Physics, Faculty of Science, Palacky University in Olomouc, 17. Listopadu 12, 77900 Olomouc, Czech Republic; [email protected] 5Institute of Physics of Materials, Czech Academy of Sciences, Zizkova 22, 61662 Brno, Czech Republic; [email protected] *Correspondence: [email protected].cz; Tel.: +420-773-782-177 Received: 21 October 2019; Accepted: 1 December 2019; Published: 5 December 2019   Abstract: Advanced thermomechanical hot rolling is becoming a widely used technology for the production of fine-grained spring steel. Different rapid phase transformations during the inductive heat treatment of such steel causes the inhomogeneous mixture of martensitic, bainitic, and austenitic phases that affects the service properties of the steel. An important task is to assess the amount of retained austenite and its distribution over the cross-section of the inductive quenched and tempered wire in order to evaluate the mechanical properties of the material. Three different analytical methods were used for the comparative quantitative assessment of the amount of retained austenite in both the core and rim areas of the sample cross-section: neutron diffraction—for the bulk of the material, Mössbauer spectroscopy—for measurement in a surface layer, and the metallographic investigations carried by the EBSD. The methods confirmed the excessive amount of retained austenite in the core area that could negatively affect the plasticity of the material. Keywords: spring steel; heat treatment; retained austenite; Mössbauer spectroscopy; neutron diffraction 1. Introduction The modern production of environmentally friendly vehicles has imparted progress in the automotive industry, leading to the development of many innovative features. The main impacts are fuel savings, which can be provided by reducing the weight of vehicles, for example, by replacing the contemporary wiring with innovative copper-aluminum clad composite wires [ 1 , 2 ], by using modern lightweight construction materials [3], or by introducing electric automobiles [4]. Due to the electric car strategy, new requirements are also given to the chassis components in terms of higher strength (above 2100 N/mm 2 ) and sufficient toughness. Several approaches could be taken in order to obtain the required results. For example, an application of the optimized treatment with the implementation of anewprocessing technology, whichcanpreferablybe done viathemethods for applying severe plastic deformation (SPD), such as equal channel angular pressing (ECAP) [5,6], accumulative roll bonding (ARB) [ 7 ], or rotary swaging [ 8 , 9 ]. The advanced treatments contribute to the optimization of the final mechanical properties via mechanical mixing and substantial grain refinement [10]. Materials 2019,12, 4063; doi:10.3390/ma12244063 www.mdpi.com/journal/materials Materials 2019,12, 4063 2 of 17 Another approach is to develop thermo-mechanically rolled fine-grained steel with a mixed microstructure consisting of martensite and/or bainite and a considerable amount of stabilized retained austenite (RA) [ 11 , 12 ]. In the last few years, many reports about the development of quenching and partitioning (Q&P) processes for the production of heat-treated steel with the optimal combination of high strength and ductility [ 13 – 19 ] can be found. All studies emphasize the key role of RA and its morphology, size, and distribution. The stability of RA also plays a key role in the prevention of tempered martensite embrittlement caused by the decomposition of RA to cementite and ferrite [20]. Grain refinement of steel microstructures was also reported to improve the stability of retained austenite by decreasing the size of blocks of RA [ 21 ]. Thermo-mechanical rolling is used to obtain the fine grain structure of the hot rolled wire products [ 22 ], along with the vanadium and/or niobium micro alloying of steel [ 23 , 24 ]. Vanadium micro-alloying also suppresses the growth of austenite grains under higher temperatures up to 1000 ◦ C during austenitization in comparison with V-free spring steel [25–27], which contributes to higher strength of the steel after heat treatment. Based on the above mentioned, the vanadium micro-alloying spring steel is supposed to provide an excellent combination of strength and ductility under optimal parameters of heat treatment. At the same time, the amount of RA and its distribution in the structure should be carefully studied because RA could be a key factor in the production of the heat-treated steel with advanced properties. Hence, the precise and reliable methods of RA assessment should be implemented. 2. Experimental Material and Methods 2.1. Material A wire rod of vanadium micro-alloyed 0.6C-0.6Mn-1.4Si-0.6Cr (wt.%) spring steel was investigated in this study. The chemical composition of the steel is presented in Table 1. A wire rod with a diameter of 17.00 mm was produced by thermo-mechanical hot rolling and was cold drawn to a diameter of 15.50 mm preceding inductive heat treatment. The microstructure of hot rolled wire rod consisted of a fine grain perlite-ferrite structure with a cementite lamellae of 17 nm and former austenite grain size of G13.0 (an average grain diameter of 3.3 µ m) according to ASTM E112 (Figure 1). The tensile mechanical properties of hot rolled wire were: ultimate tensile strength (UTS) 1350 N/mm 2 with a reduction of area value of (Z) 46%. Table 1. Chemical composition of the steel (in wt.%). Element C Mn Si Cr V Ni Mo Fe 0.560 0.580 1.400 0.570 0.150 0.024 0.002 balanced Figure 1. Austenite grain size and microstructure of hot rolled wire: ( a ) austenite grain size G13.0 according to ASTM E112; (b) cementite lamellae (transmission electron microscopy). Materials 2019,12, 4063 3 of 17 Parametersoftheinduction heat treatment are presentedinTable 2. All specimenswereinductively heated up to an austenitization temperature 850 ◦ C. Then specimens were quenched by water spraying (Q1–Q4 specimens) to the specified temperatures. Specimens Q1 and Q2 were quenched to 40 ◦ C and specimens Q3 and Q4 were quenched 100 ◦ C below calculated martensite start temperature MS=280 ◦C (Equation (1)) [ 28 ]. Temperature M f was estimated based on [ 29 ] as − 50 ◦ C with respect to the carbon content in the steel. Table 2. Parameters of the induction heat treatment. Specimen Austenitization Temperature, ◦C Temperature after Quenching, ◦C Tempering Temperature, ◦C Q1 850 40 - Q2 850 40 - Q3 850 180 - Q4 850 180 - QT1 850 40 460 QT2 850 40 420 QT3 850 180 460 QT4 850 180 300 Finally, specimens were inductively tempered at three different temperatures (Table 2) and cooled by water spraying until the ambient temperature (QT1–QT4 specimens). Tempering temperatures 460 ◦ C and 300 ◦ C were chosen as the limit values for the heat treatment of silicon alloying steel grades [ 20 ], and tempering temperature 420 ◦ C was set in order to improve the strength characteristics of the heat-treated spring steel. It should be additionally mentioned that measurements of quenched specimens Q3 and Q4 were considered as informative ones. Structural characteristics of these samples were affected by the slow air cooling until ambient temperature after the interrupted quenching by water spraying to 180 ◦ C before the measurements, as the direct investigation under 180 ◦C required comprehensive in situ study. At the same time, measurements of QT3 and QT4 specimens fully corresponded to the investigated heat treatment parameters as they were tempered immediately after reducing to 180 ◦C. MS(◦C)=521 −353C−225Si −24.3Mn −27.4Ni −17.7Cr −25.8Mo (1) Mechanical properties of the specimens were measured by means of the tensile test machine WPM UPC 1200 for quenched and tempered conditions. Hardness maps for all specimens were obtained by means of semi-automatic micro hardness equipment Durascan (Struers, Ballerup, Denmark), method HV1. Then, three different analytical methods were used in order to characterize the obtained microstructures of specimens after heat treatment. The main focus was given to the amount of RA, its distribution, and its effect on the mechanical properties of steel. 2.2. Neutron Diffraction Thefirstmethodwastheneutronpowderdiffraction(ND),whichprovidedtheaverageinformation from the whole bulk of the material. The room temperature diffraction patterns were collected on the MEREDIT instrument (Nuclear Physics Institute, ˇ Rež, Czech Republic) [ 30 ] at the Nuclear Physics Institute in Rez near Prague [ 31 ]. The mosaic copper monochromator (planes (220)) (Nuclear Physics Institute, ˇ Rež, Czech Republic) was used to provide neutrons with a wavelength of 1.46 Å. Data were collected between 4 ◦ and 144 ◦ of 2 θ with a step of 0.08 ◦ 2 θ . The samples were rotated along the vertical axis during measurement to minimize the influence of the texture and the preferred orientation on the phase fraction analysis. The full pattern structural refinements were performed using the program FullProf (Version 6.30 - Sep2018, The FullProf Suite, France) [ 32 ]. For the microstructural analysis, the instrument profile was obtained by fitting the diffraction pattern of the standard SiO 2 powder sample collected at the same conditions. Two sets of samples were investigated. The first set contained Materials 2019,12, 4063 4 of 17 the initial cylinders of as-prepared heat-treated specimens with a diameter of 15.5 mm. The second set was manufactured from the first one after the ND measurements by removing the rim part. Only the core of the original cylindrical specimens with a square profile and dimensions of 7.5 × 7.5 mm 2 was used. The 15 mm length of the specimens was kept constant for all the measurements. 2.3. Mössbauer Spectroscopy Mössbauer spectroscopy is a nuclear resonance spectroscopic technique based on the physical phenomenon of recoilless nuclear emission and resonant absorption of gamma rays [ 33 ]. This experimental technique provides qualitative and quantitative analysis of materials (e.g., structural, phase, and magnetic information) containing specific elements. The 57 Fe isotope shows the most favorable parameters for Mössbauer spectroscopy. The backscattering geometry allows us to analyze surfaces of bulk materials. Hence, 57 Fe Mössbauer spectroscopy has become a very important experimental method in steel characterization. In this study, scattering method utilized the conversion X-rays registration (the conversion X-rays Mössbauer spectroscopy—CXMS), which analyzes material surface up to depths of 1–20 µm. Nuclear hyperfine interactions can be analyzed by means of Mössbauer spectroscopy [ 34 ]. The hyperfine parameters are isomer shift, quadrupole splitting, and hyperfine magnetic splitting. The isomer shift is a result of the Coulomb interaction between the nuclear/nuclei charge and the electron charge [ 35 ]. The charges distributed asymmetrically around the atomic nucleus (electrons, ions, and dipoles) increase the electric field gradient, which differs from zero on the site of nucleus. These electric quadrupole interactions cause a splitting of the excited nuclear level [ 35 ] and provide the information about bond properties and the local symmetry of iron site. [ 33 , 34 , 36 ]. The third hyperfine parameter is magnetic splitting. This magnetic field can originate within the atom itself, within crystals via exchange interactions, or it can be external one. The magnetic field (nuclear Zeeman effect) splits the nuclear states [35]. The iron-phases analysis of the specimens was established by means of Mössbauer Spectroscopy [ 37 ]. Polished non-etched metallographic cross-sections were prepared for all specimens. The measurements were performed in a backscattering geometry at room temperature with a 57 Co(Rh) source, and spectra were recorded up to 512 channels. The isomer shift referred to the calibration of the alpha-iron sample at room temperature. The Mössbauer spectra fitting procedure was carried out using MossWinn (version 4.0, Author Dr. Zolt á n Klencs á r, Budapest, Hungary) [ 38 ] software. Paramagnetic, γ -Fe, and RA phases were fitted by one singlet. Magnetically ordered phases (ferrites, α -Fe, and those in the martensitic structure) represented by sextet components were fitted with a magnetic hyperfine field distribution [37]. The fitting process was set as free for all parameters of the Mössbauer spectra. 2.4. Electron Backscatter Diffraction (EBSD) Analysis The same cross-sections then were additionally used for EBSD analysis of the specimens Q1–Q4 in order to obtain the information about phase distribution within the rim and core areas of the specimens. Samples for EBSD were mechanically ground using SiC papers up to 4000 grit, and were polished using colloidal silica with a 0.06 µ m particle size. Finally, the specimens were electropolished using 900 mL acetic acid and 100 mL perchloric acid at 20 ◦ C for 10 s in order to remove the strains induced by mechanical preparation. The microstructures were investigated by scanning electron microscope Tescan Lyra 3 equipped with a NordlysNano detector operating at an accelerating voltage of 20 kV with the specimen tilted at 70 ◦ . The EBSD data were analyzed using HKL Channel 5 (version 5.11.10405.0, Oxford Instruments plc, Abingdon, United Kingdom) software developed by Oxford Instruments. Materials 2019,12, 4063 5 of 17 3. Results and Discussion 3.1. Mechanical Properties In order to visualize the inhomogeneity of the microstructure after fast induction heating, the hardness maps were measured for all specimens (Figure 2). Average hardness values for all specimens Q1–Q4 and QT1–QT4, along with the results of the tensile test, are shown in Table 3. Table 3. Mechanical properties of the specimens. Specimen Ultimate Tensile Strength, N/mm2Reduction of Area, % Average Hardness Value, Method HV1 Q1 - - 864 Q2 - - 862 Q3 - - 780 Q4 - - 788 QT1 2114 35 635 QT2 2176 19 653 QT3 1815 36 580 QT4 1699 0 692 Based on the results of the tensile test, only samples QT1 and QT3 had the cup and cone fracture pattern after the tensile test (Figure 3a,c) and could be used for the production of chassis components. For specimen QT2 (Figure 3b), it had a non-round cup and a cone pattern, and specimen QT4 had a brittle fracture pattern (Figure 3d), which indicates the low plasticity of these specimens; the suggested schemes of the heat treatment (QT2 and QT4) with lower tempering temperatures are not appropriate for the production of the chassis components with advanced properties. 3.2. Neutron Diffraction The neutron powder diffraction technique has an advantage for when the necessity of the average structural and microstructural information from a large volume is needed. All collected neutron diffraction patterns showed the presence of two crystallographic phases. An example of a measured and calculated neutron diffraction pattern of the specimen QT3 is presented in Figure 4. The minor phase reflections can be indexed as a face-centered cubic (FCC) lattice with a cell parameter of about 3.58 Å. This phase was recognized as RA. The major phase reflections could be within the first approximation indexed as a body-centered cubic (BCC) lattice with a cell parameter of about 2.87 Å, which looked like a fingerprint of the ferrite phase. From the crystallographic point of view, there was no possibility to distinguish a difference between ferrite and low-carbon α -martensite phase. α -martensite had the same lattice as ferrite, but due to the high dislocation and defect density, the strong microstructural reflection broadening or asymmetry could be observed, but the quantitative evaluation of the phase ration was not possible. The detailed analysis of the neutron diffraction profile showed the significant reflection broadening ofbothphasessurpassingtheinstrumentbroadening(significantspecimencontribution). Itindicatesthe presence of an increased microstrain in the specimens. In addition, in the case of the specimens Q1–Q4, the shape of the reflections of the ferrite phase was found to be strongly asymmetric, indicating the split of the cell parameters and the decrease of the symmetry from cubic to tetragonal. The appearance of the tetragonality indicates the presence of the martensitic phase. The results of the structural and microstructural analysis are depicted below. Materials 2019,12, 4063 6 of 17 Figure 2. Hardness mapping of the specimens, method HV1: ( a ) Q1; ( b ) QT1; ( c ) Q2; ( d ) QT2; ( e ) Q3; (f) QT3; (g) Q4; (h) QT4. Materials 2019,12, 4063 7 of 17 Figure 3. Fracture surface of the QT specimens: ( a ) QT1—cup and cone pattern; ( b ) QT2—non-round cup and cone pattern; (c) QT3—cup and cone pattern; (d) QT4—brittle fracture. Figure 4. Measured (red crosses) and calculated (black line) neutron diffraction pattern of the specimen QT3 together with their difference (blue bottom line). Small vertical bars represent the Bragg positions for ferrite and austenite, respectively. The amount of RA calculated from the full pattern fitting is presented in Figure 5as a function of quenching and tempering temperatures, i.e., parameters of the heat treatment. The comparison of the measurements of the whole specimen and the core area confirmed that the excessive amount of Materials 2019,12, 4063 8 of 17 RA was located at the specimen center. The reason for the higher amount of RA in the core area was possibly the relatively slow heat transfer from the surface to the core during the quenching process, which leads to the significantly slower and delayed transformation within the specimen core. Figure 5. Weight fraction of retained austenite in (%). In order to evaluate the influence of RA on the mechanical properties of the material, specimens QT1 and QT3 should be compared. The amount of RA in the bulk of the specimen QT1 is 2.4 times lower than in QT3, but both samples showed the reduction of area ≥ 35%, which is acceptable for the production of chassis components. It indicates that ductility is not affected by the higher amount of stabilized RA in the QT3 specimen, and it reaches similar values for both QT1 and QT3 specimens. Contrarily, UTS and hardness show a significant decrease for the QT3 specimen with the higher amount of the soft FCC (RA) phase. The change in lattice parameter for ferrite/ α -martensite and austenite, which perfectly describes its evolution with respect to the different parameters of heat treatment, can be seen in Figure 6. As the speed of induction heating was considered to be constant for all samples, it can be concluded that the evolution of the lattice parameter was attributed to the different temperatures and, therefore, to the carbon activity and its diffusion under such temperatures. The highest value of tetragonality, i.e., the c/a parameter, was observed for quenched specimens Q1 and Q2 (1.008 in the bulk), which confirmed the high amount of martensitic phase in the structure (Figure 6) [ 39 ]. On the contrary, the tetragonality of all QT specimens was zero, which suggests that the tempering temperature was sufficient for carbon diffusion from BCC lattice. The lattice parameter for the FCC phase was the opposite of the BCC evolution. Obtained results also correspond with the description of phase transformation in the TRIP C-Mn-Si sheet steel studied by Yu et al. [ 40 ] by means of in situ neutron diffraction. Yu et al. attributed the expansion of the lattice parameter of FCC to the carbon enrichment of RA, and expected the shrinkage of FCC lattice after further tempering above 470 ◦C due to diffusion of the carbon from RA to ferrite. Materials 2019,12, 4063 9 of 17 Figure 6. Lattice parameters: (a) body-centered cubic – ferrite; (b) face-centered cubic – austenite. The evaluation of the reflection profile using an anisotropic strain-broadening model incorporated in the FullProf program (Version 6.30 - Sep2018, The FullProf Suite, France) revealed an average and directional µ -strain (combination of strain of type-II and type-III) of the individual phases. Figures 7and 8 show the values of the µ -strain for all specimens, both for the bulk and the core areas for the ferrite and austenite phase, respectively. A strong directional asymmetry was found for the ferrite phase ( µ -strain along 100 is significantly higher than the one along 111, see Figure 7) in comparison with the austenite phase where the asymmetry is minimal (represented by error bars in Figure 8). Materials 2019,12, 4063 16 of 17 15. Gao, G.; Zhang, B.; Cheng, C.; Zhao, P.; Zhang, H.; Bai, B. 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